Method of sensing a gas by detecting change in magnetic properties

ABSTRACT

An oxide semiconductor doped with a transition metal and exhibiting room-temperature ferromagnetism is disclosed. The transition metal-doped oxide semiconductor is preferably manufactured in powder form, and the transition metal is preferably evenly distributed throughout the oxide semiconductor. The preferred embodiments are iron-doped tin dioxide and cobalt-doped tin dioxide. Gases may be detected by passing them across a material and measuring the change in magnetic properties of the material; the preferred material is iron-doped tin dioxide.

This application claims priority based on U.S. Provisional ApplicationNo. 60/598,203, Ferromagnetic Powders of Tin Oxide Nanoparticles Dopedwith Cobalt and Magnetic Gas Sensor Utilizing Them, filed Jul. 30, 2004,and U.S. Provisional Application No. 60/612,708, Development of HighTemperature Ferromagnetism in SnO₂ and Paramagnetism in SnO by FeDoping, filed Sep. 23, 2004, the disclosures of both of which are herebyincorporated by reference.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to gas sensors which detect gas bymeasuring the change in magnetic properties of a ferromagnetic material.

2. Related Art

Solid materials are either crystalline or amorphous. A crystalline solidis one in which the atomic arrangement is regularly repeated, and whichis likely to exhibit an external morphology of planes makingcharacteristic angles with each other. In many materials, there areactually a variety of solid phases, each corresponding to a uniquecrystal structure. These varying crystal phases of the same substanceare called “allotropes” or “polymorphs”. The mechanical, thermal,optical, electronic, and magnetic properties of crystals are stronglyinfluenced by the periodic arrangement of their atomic cores. Ananoscale particle is a particle having a measurement of 100 nm or lessin at least one direction.

A crystal may be regarded as a three-dimensional diffraction grating forenergetic electromagnetic waves (typically X-rays) of a wavelengthcomparable with the atomic spacing; the diffraction pattern will provideinformation about the periodic arrangements of the atoms. Constructiveinterference of the electromagnetic waves may occur where the followingminimum condition (called the Bragg equation) is satisfied:2d sin θ=nλwhere d is the spacing between crystalline planes, θ is the angle ofincidence between the beam of X-rays and the parallel crystallineplanes, n is an integer, and λ is the wavelength of the X-rays. Thisequation will not be satisfied for most angles θ; to investigatecrystals of unknown orientation, the rotating crystal method is used,wherein θ is varied as a function of time while X-rays of a singlewavelength are presented to a single crystal. A cylinder of photographicfilm records a spot whenever the Bragg condition is fulfilled.

In the Debye-Scherrer technique, instead of using a single X-raywavelength and a time-dependent angle of incidence, a crystalline sampleis presented with every θ simultaneously. This is achieved by using afinely powdered crystalline sample in which the crystalline orientationsare random. Rays for which one crystallite or another satisfy the Braggcondition emerge from the sample as a series of cones concentric withthe incident beam direction. Thus a photographic plate records a seriesof concentric circles. The spacing and pattern of these circles is usedto determine the atomic structure of the crystal.

Particle-induced x-ray emission (PIXE) is an analytical techniquecapable of trace element detection sensitivity of a few parts permillion. When ions pass through matter, they interact with the electronsin the atoms and occasionally a vacancy is produced by an excitedelectron. When this occurs in an inner shell, the vacancy is filled byan electron from an outer shell, and an x-ray photon of characteristicenergy is emitted. By measuring the energy of the photon, one candetermine the atomic number of the element and the amount of the elementpresent that can be extracted from the area under the x-ray peak. Foridentification and quantification of trace elements, PIXE is 100 timesmore sensitive than electron micro-analysis systems.

Mossbauer spectroscopy is a spectroscopic technique based on theMossbauer effect. In its most common form, Mossbauer AbsorptionSpectroscopy, a solid sample is exposed to a beam of gamma radiation,and a detector measures the intensity of the beam that is transmittedthrough the sample. The gamma-ray energy is varied by accelerating thegamma-ray source through a range of velocities with a linear motor. Therelative motion between the source and the sample results in an energyshift due to the Doppler effect. In the resulting spectra, gamma-rayintensity is plotted as a function of the source velocity. At velocitiesresponding to the resonant energy levels of the sample, some of thegamma-rays are absorbed, resulting in a dip in the measured intensityand a corresponding dip in the spectrum. The number, positions, andintensities of the dips (also called peaks) provide information aboutthe chemical environment of the absorbing nuclei and can be used tocharacterize the sample.

In x-ray photoelectron spectroscopy, the sample is illuminated with softx-radiation in an ultrahigh vacuum. The photoelectric effect leads tothe production of photoelectrons, the energy spectrum of which can bedetermined in a beta-ray spectrometer. The difference between the x-rayphoton energy, which is known, and the electron energy, which can bemeasured, results in the binding energy of the orbital from which theelectron was expelled. Measurement of the relative areas of thephotoelectron peaks allows the composition of the sample to bedetermined.

In discussing magnetic fields, the relationship between the magneticfield intensity H and the corresponding magnetic induction B is$\begin{matrix}{B = {µ_{0}\left( {H + M} \right)}} \\{= {{µ_{0}\left( {1 + \chi_{m}} \right)}H}} \\{= {µ_{0}µ_{m}H}} \\{= {µ\quad H}}\end{matrix}$where M is the magnetization, χ_(m) is the magnetic susceptibility,μ_(m) is the relative permeability, μ is the absolute permeability, andμ₀=4π×10⁻⁷H/m is the permeability of free space.

The magnetic response of most solids is dominated by the orientation ofpermanent dipoles. The response of a magnetic material is usuallyexpressed in terms of either the magnetization M or the magneticsusceptibility χ, whereM=χHandχ=M/H

A spinning charged particle constitutes a magnetic dipole. The magneticdipole moment of an electron is attributed to its “spin,” and creates amagnetic field pointing in a direction perpendicular to the plane inwhich the electron is spinning, as shown in FIG. 1.

There are four different kinds of magnetic behavior which involvepermanent dipoles in a solid, namely paramagnetic, antiferromagnetic,ferromagnetic, and ferrimagnetic. The low temperature ordering, if any,of neighboring dipoles, and the consequent behavior of spontaneousmagnetization and/or susceptibility results in hysteresis loops (shownin FIG. 3) in ferromagnetic and ferromagnetic materials.

Paramagnetic behavior, shown in FIG. 2(a), occurs when the magneticmoments of the various atoms are uncorrelated in the absence of amagnetic field, and the sum total of the magnetic moments tends towardzero. The dipoles do tend to become aligned in a magnetic field. Themagnetic susceptibility follows the Curie Law:χ_(m) =C/Twhere C is the Curie constant of the solid, and T represents thetemperature of the solid.

Antiferromagnetic behavior, shown in FIG. 2(b), occurs when the dipoles,or magnetic moments, alternate, causing the sum total of the magneticmoments to tend toward zero. This arrangement is very stable at lowtemperatures, and the magnetic susceptibility in an applied field issmall. When the temperature rises, the efficiency of this dipole-dipoleinteraction decreases and the magnetic susceptibility increases, untilthe spins become “free” at the Neel temperature to respond to a field.At even higher temperatures the behavior becomes paramagnetic, and themagnetic susceptibility follows a modified Curie lawχ_(m=) C/(T+θ)

A ferromagnetic solid, represented in FIG. 2(c), is ordered withparallel spins below the Curie temperature T_(c), which results in aspontaneous magnetization M_(s). The magnitude of this bulk polarizationdecreases to zero at the Curie temperature T_(c) (which is well belowroom-temperature in most ferromagnetic solids), and the paramagneticsusceptibility for the disordered spin system at higher temperaturesobeys the Curie-Weiss lawχ_(m) =C/(T−T _(c))

Ferromagnetism involves the cooperative alignment of permanent atomicdipoles, which arise in atoms having unpaired electrons. The strength ofeach individual dipole is small, but a completely ordered array of suchmoments produces a large spontaneous magnetization M_(s).

The low temperature ordering in a ferrimagnetic material, as shown inFIG. 2(d), is similar to that of an antiferromagnetic material, but thetwo opposing spin systems have magnetic moments of unequal magnitude,and a net spontaneous magnetization results. This magnetization declinesto zero magnitude when the solid is warmed to the Curie point T_(f), andthe behavior is once again paramagnetic at higher temperatures.

Hysteresis loops demonstrate a phenomenon wherein a material that didnot show any magnetization before the application of a magnetic fieldexhibits remanent magnetization after the applied magnetic field isremoved, as shown in FIG. 3. The coercivity H_(c) is the value of themagnetic field that must be applied to return the magnetization to zeroafter the magnetization has been caused to reach its saturation value;the remanence B_(r) (M_(r) in the text below) is the value of themagnetization of the material after the material has been caused toreach its saturation value and then had the applied magnetic fieldremoved. Non-zero values for the coercivity and remanence of a sampleimply that the sample is ferromagnetic.

Oxide semiconductors have been used to detect gases. However, these haveall been non-magnetic oxide semiconductors. Their electrical andsemiconducting properties (determined by electrical resistivity, carrierconcentration, and carrier mobibility) vary with oxygen stoichiometry.Oxygen stoichiometry can be changed by passing a reducing or oxidizinggas. Thus, traditionally, monitoring the changes in the electricalproperties with the type and flow rates of gases has been used as asensing method.

Tin Dioxide, SnO₂, is an oxide semiconductor with a wide band gap of˜3.6 eV. When prepared with oxygen vacancies, SnO₂ becomes an n-typesemiconductor. When doped with iron (Fe) or cobalt (Co), SnO₂ remains ann-type semiconductor. Development of room-temperature ferromagnetism(RTFM) in conventional semiconductors is currently attracting intenseinterest due to their potential use in spintronics applications.However, most traditional transition metal-doped magnetic semiconductorsystems exhibit ferromagnetism only at temperatures that are well belowroom temperature.

SUMMARY OF THE INVENTION

The present invention is a gas sensor that detects the presence of a gasby measuring the change in the magnetic properties of a ferromagneticmaterial as the gas flows across the material. This detector has beenenabled by the invention of a material exhibiting ferromagnetism at roomtemperature, preferably in powder form. The preferred material isiron-doped tin dioxide.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows a ring of charge rotating about its axis.

FIG. 2 shows the low temperature ordering, if any, of neighboringdipoles, and the consequent behavior of spontaneous magnetism and/orsusceptibility, for (a), paramagnetism, (b) antiferromagnetism, (c)ferromagnetism, and (d) ferrimagnetism.

FIG. 3 shows a schematic magnetization loop for a multi-domain sample ofa ferromagnetic solid. H_(c) is the coercivity, and B_(r) is theremanence of the sample. The dashed curve shows what happens when anonmagnetized sample is first magnetized. The arrows on the solid curvesshow the course of a subsequent hysteresis loop.

FIG. 4 shows typical PIXE spectra from the Sn_(1-x)Fe_(x)O₂ samplesshowing the Fe region. Fe concentrations (x) obtained by PIXE datasimulation (solid lines) are given in parentheses.

FIG. 5 shows (a) typical PIXE spectra from the Sn_(1-x)Co_(x)O₂ samplesshowing the Co region, with Co concentrations (x) obtained by datasimulation (solid lines) also included; (b) XRD data of Sn_(1-x)Co_(x)O₂powders prepared with different values of x; powder diffraction files oforthorhombic SnO₂ (solid thin lines, marked O), tetragonal SnO₂ (solidthick lines, marked C), and SnO (thick dashed lines, marked R) phasesare also shown; the * indicates weak peaks of Co₃O₄ observed only insamples with x≧0.08; (c) TEM images of Sn_(1-x)Co_(x)O₂ prepared at 600°with x=0.01; the inset shows electron diffraction patterns taken fromselected regions of the TEM images; (d) TEM images of Sn_(1-x)Co_(x)O₂prepared at 600° with x=0.05; the inset shows electron diffractionpatterns taken from selected regions of the TEM images; (e) TEM imagesof Sn_(0.99)Co_(0.01)O₂ prepared at 350° C.; (f) XRD data ofSn_(1-x)Co_(x)O₂ powders prepared with by annealing at differenttemperatures; the bottom and top panels show data from samples annealedin air and in flowing hydrogen (10% H₂ and 90% He) respectively; powderdiffraction files of orthorhombic SnO₂ (solid thin lines, marked O),tetragonal SnO₂ (solid thick lines, marked C), and SnO (thick dashedlines, marked R) phases are also shown.

FIGS. 6(a) and 6(b) show XRD patterns of Sn_(1-x)Fe_(x)O (prepared at200° C.), and Sn_(1-x)Fe_(x)O₂ (prepared at 600° C.), respectively,along with reference lines of orthorhombic SnO₂ (solid lines, marked“O”), romarchite SnO (dotted lines, marked “R”), cassiterite SnO₂(dashed lines, marked “C”) phases, hematite (marked “H”), and maghemite(marked “M”) phases of Fe₂O₃.

FIG. 7 shows (a) changes in the lattice parameters a and c of tetragonalSnO calculated using (101) and (110) peaks as a function of Fepercentage, as well as the reported magnitude of the lattice parametersof bulk SnO; (b) changes in the lattice parameters a and c ofcassiterite SnO₂ calculated using the (110) and (202) peaks as afunction of Fe percentage, as well as the reported magnitude of thelattice parameters of bulk SnO₂; (c) particle size of Sn_(1-x)Fe_(x)O₂as a function of x calculated from the tetragonal cassiterite XRD peak(110), with particle sizes determined from TEM marked with stars.

FIGS. 8(a) and (b) show (a) XRD patterns of 5% Fe-doped samples preparedby annealing the reaction precipitate at different temperatures shownabove, along with reference lines of orthorhombic SnO₂ (solid lines,marked “O”), romarchite SnO (dotted lines, marked “R”) cassiterite SnO₂(dashed lines, marked “C”) phases, hematite (marked “H”) and maghemite(marked “M”) phases of Fe₂O₃; (b) Changes in the lattice parameter a andthe lattice volume of cassiterite Sn_(0.95)Fe_(0.05)O₂ as a function ofpreparation temperature. The change in the lattice parameter c wasminimal.

FIGS. 9(a)-(d) show (a) intensity of the tetragonal cassiterite peak(110) and the orthorhombic fraction x₀ of Sn_(1-x)Co_(x)O₂ prepared at600° C. as a function of Co percentage; (b) the particle size ofSn_(1-x)Co_(x)O₂ as a function of x calculated from the XRD tetragonalcassiterite peak (110). The particle sizes determined from TEM aremarked with stars; (c) changes in the lattice parameters a and c ofcassiterite SnO₂ as a function of Co percentage. Stars indicate the bulkvalues of the SnO₂ lattice parameters from XRD reference files; (d)changes in the (110) cassiterite peak position with Co concentration.The lattice parameters were calculated using (110) and (202) peaks ofthe tetragonal cassiterite phase.

FIGS. 10(a) and (b) show diffuse reflectance spectra of Sn_(1-x)Co_(x)O₂samples prepared at 600° C. (a) The changes in the absorption edge withCo concentration; inset in panel (a) shows the changes in the band gapenergy estimated from the reflectance data as a function of Coconcentration and (b) shows the complete spectra indicating the extentof Co₃O₄ formation.

FIGS. 11(a)-(c) show (a) Raman spectra of Sn_(1-x)Co_(x)O₂ prepared at600° C. as a function of x; (b) Raman spectra of pure SnO₂, 1% Co-dopedSnO₂ and a physical mixture of pure SnO₂ and Co₃O₄; (c) the apparentdisappearance of SnO₂ Raman peak at 630 cm⁻¹ and the emergence of the617 cm⁻¹ peak of Co₃O₄ for x>0.01.

FIGS. 12(a)-(c) show panels (a) and (b) which show transmission electronmicroscopy (TEM) images of Sn_(1-x)FexO₂ prepared at 600° C. with x=0.01and 0.05, respectively. Panel (c) shows the TEM image ofSn_(1-x)Fe_(x)O₂ prepared at 200° C. with x=0.05.

FIGS. 13(a) and (b) show panels (a) and (b) which show TEM images ofSn_(0.95)Fe_(0.05)O₂ prepared at 350 and 900° C. respectively.

FIGS. 14(a)-(c) show panel (a) which shows the room-temperatureMossbauer spectra of Sn_(0.95)Fe_(0.05)O. Panels (b) and (c) show theroom-temperature Mossbauer spectra of Sn_(0.95)Fe_(0.05)O₂ prepared at350 and 600° C., respectively.

FIG. 15 shows XPS spectra of Sn_(1-x)Fe_(x)O₂ prepared at 600° C. withdifferent values of x as indicated. Reference data obtained frommaghemite and hematite forms of Fe₂O₃ prepared under identical synthesisconditions (but with no Sn precursors) are also shown.

FIG. 16 is a plot showing the XPS spectra of 5% Fe-doped samples as afunction of the annealing temperature. Reference data obtained frommaghemite and hematite forms of Fe₂O₃ prepared at 200 and 600° C.,respectively, under identical synthesis conditions (but with no Snprecursors) are also shown.

FIG. 17 is a plot showing the XPS spectra of Sn_(0.95)Fe_(0.05)O₂prepared at 900° C. along with that obtained from the same sample afterremoving 10 and 20 nm of surface layer by Ar⁺ ion sputtering.

FIGS. 18(a) and (b) show (a) XPS spectra of Sn_(1-x)Co_(x)O₂ prepared at600° C. as a function of Co percentage and (b) shows similar data ofCo₃O₄ (x=1) and SnO₂ (x=0) reference samples prepared under identicalsynthesis conditions.

FIGS. 19(a) and (b) show the variation of the atomic percentages of Co,Sn, and O and the Co/Sn ratio as a function of Co concentrationcalculated using the corresponding XPS peak intensities.

FIGS. 20(a)-(c) show (a) and (c) M vs H data of Sn_(1-x)Fe_(x)O₂measured at 300 and 5K, respectively; (b) M−χ_(P)H as a function of Hfor the Sn_(1-x)Fe_(x)O₂ samples measured at 300K. Solid lines throughthe data points in (c) are theoretical fits using the modified Brillouinfunction for a paramagnetic system.

FIG. 21 shows M vs T data measured with H=500 Oe for pure iron oxidesamples prepared at 200° C. (maghemite) and 600° C. (hematite).

FIG. 22 shows changes in the saturation magnetization Ms and remanenceM_(r), along with the XPS estimate of surface Fe concentration of theSn_(0.95)Fe_(0.05)O₂ samples as a function of their preparationtemperature.

FIGS. 23(a) and (b) show XRD data collected from separate angularregions for Sn_(0.95)Fe_(0.05)O₂ (data points), pure SnO₂ (dashed line),and pure Fe₂O₃, all prepared at 600° C. following identical synthesisprocedures. The intensity is plotted on a log scale.

FIGS. 24(a)-(c) show changes in (a) saturation magnetization M_(s) andlattice volume V, (b) remanence M_(r) and the linear paramagneticcomponent χ_(p), and (c) interaction parameter T₀ and Curie-Weisstemperature θ (obtained from the paramagnetic component in FIG. 25) ofthe Sn_(1-x)Fe_(x)O samples as a function of x.

FIG. 25 shows M vs T data measure with H=500 Oe from Sn_(1-x)Fe_(x)O₂samples. Solid lines are theoretical fits using the modified Curie-Weisslaw.

FIG. 26 shows M vs T data measured with H=500 Oe from Sn_(1-x)Fe_(x)Osamples. Solid lines are theoretical fits using the modified Curie-Weisslaw.

FIG. 27 shows temperature variations of the magnetic susceptibility X,measured with an applied field H=500 Oe, for the 1% and 3% Co dopedsamples prepared at 600° C. Lines through the data points aretheoretical fits using the modified Curie-Weiss law. The inset shows themagnetization of the same samples, measured at 5K as a function ofmagnetic field. Lines through the data points are theoretical fits usingthe Brillouin-function-based form for a paramagnetic system.

FIG. 28 shows room-temperature hysteresis loops of Sn_(1-x)Fe_(x)O₂prepared at 600° C. with x=0.05 (main panel) and 0.01 (inset). The linesjoining the points are for visual aid.

FIG. 29 is a plot of the normalized sample magnetization M of aSn_(0.99)Fe_(0.01)O₂ sample measured with H=10 kOe as a function oftemperature, indicating a Curie temperature T_(C)˜850 K.

FIG. 30 shows a room temperature hysteresis loop obtained fromSn_(0.99)Co_(0.01)O₂ prepared at 350° C. The insets (a) and (b) show thehysteresis loops of Sn_(0.99)Co_(0.01)O₂ prepared at 600° C. andSn_(0.995)Co_(0.005)O₂ prepared at 350° C. The lines joining the pointsare for visual aid.

FIGS. 31(a) and (b) show M vs H data for Sn_(1-x)Fe_(x)O samplesmeasured at 5 and 300 K, respectively. Similar M vs H data collectedfrom pure iron oxide (maghemite) prepared under identical conditions(but with no Sn precursors) are also shown. Solid lines through the datapoints are theoretical fits using the modified Brillouin-function-basedform for a paramagnetic system.

FIGS. 32(a)-(c) show (a) the low field region of the room-temperaturehysteresis loop of 600° C. prepared Sn_(0.99)Co_(0.01)O₂ sample showinga coercivity of 9 Oe; the inset in (a) shows the complete hysteresisloop of 600° C. prepared Sn_(0.99)Cu_(0.01)O₂ showing saturation of thesample magnetization expected for a ferromagnetic system; (b) variationof the room-temperature coercivity H_(c) and remanence M_(r) ofSn_(1-x)Co_(x)O₂ prepared at 350° C. as a function x; inset in (b) showsthe expanded view of the low-field region of the hysteresis loop of 350°C. prepared Sn_(0.995)Co_(0.005)O₂ sample; and, (c) variation of thesaturation magnetization Ms with x of Sn_(1-x)Co_(x)O₂ samples preparedat 350 (open circles) and 600° C. (solid stars), and the lattice volumeV calculated using the a and c values of FIG. 9(c) (shown with solidsquares) as a function of Co-doping concentration for the 600° C.prepared Sn_(1-x)Co_(x)O₂.

FIG. 33 shows M vs H data of Sn_(0.95)Co_(0.05)O₂, prepared at thedifferent temperatures indicated, measured at 300 K.

FIG. 34 shows an embodiment of a gas sensing apparatus using anembodiment of the invented process for detecting a gas.

FIG. 35 shows changes in the saturation magnetization ofSn_(0.95)Fe_(0.05)O₂ versus flow rate of molecular oxygen at 200° C. fora sample prepared at 600° C.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

1. The Materials Used in the Gas Sensing Process

The preferred embodiment is a powder comprising an oxide semiconductorthat is transition-metal doped; preferably, the semiconductor is SnO₂.Impurities from other elements could be present and not significantlyaffect the magnetic properties of the composition. The invention doesnot have to be in powder form, as the composition manufactured by thepreferred process can be converted into other forms, such as films. Inthe preferred embodiment, the transition metal is iron (Fe), the dopingconcentration is between 0.5% and 10%, the Curie temperatures is as highas 850K, the composition takes on a powder form with nanoscale particlesof which 95% are believed to be less than 100 nm in length, there are nophases, or clusters of Fe, in the composition, meaning that the Fe isevenly distributed throughout the composition, and the composition isintrinsically ferromagnetic, meaning that the Fe atoms take the place ofthe Sn atoms in the lattice, and are substitutionally incorporated intothe SnO₂ lattice at the Sn sites. Preparation conditions have a strongeffect on the observed magnetic properties and might act as a usefulcontrol parameter.

The preferred embodiment, Sn_(1-x)Fe_(x)O₂, is manufactured by thefollowing preferred process, which is less expensive than other methodsof manufacturing ferromagnetic oxide semiconductors. Appropriate amountsof tin dichloride (SnCl₂) of minimum 99% purity, iron dichloride (FeCl₂)of minimum 99.5% purity, and NH₄OH are added to de-ionized water toproduce solutions with molarities of 1, 0.02, and SM, respectively. Allthe samples are prepared by reacting the 0.02M FeCl₂ and 1 M SnCl₂solutions at 80° C. {molar ratio of x=[Fe]/([Fe]+[Sn])} with a largeamount (˜1.5 times the precursor solution volume) of a SM solution ofNH₄OH. The resulting precipitate is washed to remove any water-solublebyproducts and annealed in air for three hours at 600° C. to obtainpowder samples of Sn_(1-x)Fe_(x)O₂; in the case of the ratios of 0.02 MFeCl₂ and 1 M SnCl₂ in one embodiment, Sn_(0.98)Fe_(0.02)O₂ is obtained.Samples of Sn_(0.95)Fe_(0.05)O₂ have also been prepared by annealing thesame precipitate at temperatures of 350, 450, 750, and 900° C. for thepurpose of investigating the effect of the annealing temperature. Whenthe precipitate is annealed at 200° C., iron-doped tin monoxide(Sn_(1-x)Fe_(x)O) results. Pure iron oxide samples have been preparedfollowing identical synthesis procedures without using any SnCl₂ toobtain insight into possible Fe impurity phases that might form underthese synthesis conditions.

In an alternative embodiment, cobalt-doped tin dioxide, Sn_(1-x)Co₂O₂,with x being 1% or less, exhibiting room-temperature ferromagnetism, hasbeen developed. In this alternative embodiment, magnetic hysteresisloops are observed at 300 K (room-temperature) with coercivity Hc ˜630Oe, saturation magnetization M_(s)˜0.233 μ_(B)/Co ion and about 31%remenance. SnO₂ samples doped with ≦1% Co showed RTFM with significantlyhigh coercivity (˜630 Oe), moderate remenance (˜31%) and bettersquareness of the hysteresis loop, but had a lower magnetic moment of0.133 μ_(B)/Co ion. However, for x>0.01, this ferromagnetism wascompletely destroyed and the samples demonstrated paramagnetic behavior.Preferably, the Co is evenly distributed throughout the SnO₂ lattice,and the composite takes on a nanoscale particle form.

The Sn_(1-x)Co_(x)O₂ is preferably prepared using a wet chemical methodby reacting 0.02 M CoCl₂.6H₂O and 1 M SnCl₂ with a molar ratio ofx=[Co/(Co+Sn)]. A few drops of concentrated HCl are added to ensuredissolution. This solution is added to a 5M solution of NH₄OH, and theresulting mixture is heated to 80° C. for several hours. The precipitateis annealed in air at various temperatures for three hours to obtainSn_(1-x)Co_(x)O₂. Chemically synthesized Sn_(1-x)Co_(x)O₂ powders havebeen shown to exhibit RTFM for x≦0.01 when prepared in the 350 to 600°C. range.

The sol-gel based wet chemistry used to manufacture the two preferredembodiments is preferred over other possible means because it isrelatively inexpensive, it intrinsically excludes the segregation oftransition metal nanoparticles, it has the ability to kineticallystabilize metastable phases (such as orthorhombic SnO₂) and extendedsolid solutions, and it is very efficient for the controlled synthesesof materials in the nanosize range.

For both Fe- and Co-doped SnO₂ with preparation temperatures of <300° C.and >750° C., there is no ferromagnetism and the particles are notnanoscale in size. It is believed that a ferromagnetic powder withnanoscale particles can be achieved with an annealing temperature above300° C. and below 750° C.

A. Confirmation of Nominal Doping Concentrations

The nominal Fe doping concentrations of both Sn_(1-x)Fe_(x)O₂ andSn_(1-x)Fe_(x)O have been confirmed by PIXE measurements. The powdersamples were first mixed with a very small amount of polyvinyl alcoholand then palletized using a hand-held press. The samples were thenirradiated with a 2.0 MeV He+ ion beam and the x-rays emitted during thede-excitation process within the atoms were analyzed using an x-rayspectrometer.

The PIXE data obtained from selected samples of Sn_(1-x)FexO₂ are shownin FIG. 4. The Fe concentrations shown in Table 1, estimated bysimulating the experimental PIXE spectra after removing the backgrounddue to bremsstrahlung, are in reasonable agreement with their nominalconcentrations. This confirms that the ratio of Fe to Sn in theresulting product is substantially the same as the ratio of the FeCl₂and SnCl₂ solutions that were reacted with NH₄OH. TABLE I Atomicconcentration estimates of Sn_(1−x)Fe_(x)O and Sn_(1−x)Fe_(x)O₂ obtainedfrom PIXE and XPS measurements. Prepa- ration Nom- Fe % tem- Major XRDEstimated atomic inal from perature identified Processing % from XPS Fe% PIXE (° C.) phase conditions Fe Sn O 0 — 200 Sn_(1−x)Fe_(x)OAs-prepared — 29.2 55.4 1 0.79 200 Sn_(1−x)Fe_(x)O As-prepared — 30.161.0 3 3.07 200 Sn_(1−x)Fe_(x)O As-prepared — — — 5 4.54 200Sn_(1−x)Fe_(x)O As-prepared 1.2 29.2 61.0 0 — 600 Sn_(1−x)Fe_(x)O₂As-prepared — 31.2 61.2 1 0.66 600 Sn_(1−x)Fe_(x)O₂ As-prepared — 31.762.2 3 2.66 600 Sn_(1−x)Fe_(x)O₂ As-prepared — — — 5  4.888 600Sn_(1−x)Fe_(x)O₂ As-prepared 3.1 27.7 59.7 5 — 350 Sn_(1−x)Fe_(x)O₂As-prepared 0.7 30.0 60.2 5 — 450 Sn_(1−x)Fe_(x)O₂ As-prepared 1.1 31.864.2 5 — 750 Sn_(1−x)Fe_(x)O₂ As-prepared 3.9 28.6 62.3 5 4.00 900Sn_(1−x)Fe_(x)O₂ As-prepared 6.2 23.1 56.3 5 — 900 Sn_(1−x)Fe_(x)O₂ 10nm Ar⁺ 4.5 36.3 57.5 ion sputtered 5 — 900 Sn_(1−x)Fe_(x)O₂ 20 nm Ar⁺3.4 39.7 56.2 ion sputtered

Typical PIXE spectra from Sn_(1-x)Co_(x)O₂ samples are shown in FIG.5(a). The Co concentrations of 1.4, 3.5 and 5.4% estimated by simulatingthe experimental PIXE spectra after removing the background due tobremsstrahlung are in reasonable agreement with their nominalconcentrations of 1, 3, 5% Co respectively.

B. Crystalline Structure of the Compositions

X-ray diffraction (XRD) studies utilizing the Debye-Scherrer techniquewere used to determine the crystalline structure of the compositionsobtained. XRD spectra were recorded at room temperature on a PhillipsX′Pert x-ray diffractometer with a CuK_(α) source (μ=1.5418 Å) isBragg-Brentano geometry. The loose powder samples were leveled in thesample holder to ensure a smooth surface and mounted on a fixedhorizontal sample plane. Data analyses were carried out using profilefits of selected XRD peaks.

As shown in FIG. 6(b), the powder Sn_(1-x)Fe_(x)O₂ samples showed strongXRD peaks due to the cassiterite phase of SnO₂, with much weaker peaksof the metastable orthorhombic SnO₂ phase. The peak intensities,positions, and widths of the XRD lines changed with x inSn_(1-x)Fe_(x)O₂, as shown in FIG. 6(b). Lattice parameters a and c andthe particle size L, as shown in FIGS. 7(b) and 7(c), estimated usingthe cassiterite (110) and (202) peaks of the nanoscale samples ofSn_(1-x)Fe_(x)O₂, decreased as x increased from 0.005 to 0.05.

The XRD patterns of powder Sn_(1-x)Fe_(x)O samples, on the other hand,showed strong peaks of tetragonal SnO with some weak SnO₂ traces, asshown in FIG. 6(a). The lattice parameters a and c, determined using the(101) and (110) peaks, showed an increase with x, as shown in FIG. 7(a).The experimentally determined lattice parameters of the pure SnO samplesare lower than that reported for pure synthetic bulk romarchite; thismay be due to changes in the oxygen stoichiometry and/or particle sizeeffect.

The directly opposite changes in the lattice parameters observed inSn_(1-x)Fe_(x)O₂ and Sn_(1-x)Fe_(x)O with Fe doping concentration mightreflect the effect of substituting Fe³⁺ for Sn⁴⁺ ions in SnO₂ and forSn²⁺ ions in SnO. This might require rearrangement of neighboring oxygenions for charge neutrality. When the 5% Fe-doped Sn_(1-x)Fe_(x)O₂samples were prepared at different temperatures in the 200 to 900° C.range, the tetragonal SnO phase was observed at 200° C. and showed agradual conversion to the SnO₂ phase with increasing preparationtemperature until its apparent disappearance at ≧450° C., as illustratedin FIG. 8(a). The lattice parameter a and the unit cell volume Vdecreased, and the lattice parameter c increased with increasedpreparation temperature in the 350 to 600° C. range, as shown in FIG.8(b). Above 600° C., these trends were reversed, and the lattice volumeapproached closer to the pure SnO₂ range.

XRD patterns of the Sn_(1-x)CO_(x)O₂ samples showed the formation oftetragonal cassiterite SnO₂ with a very small fraction of metastableorthorhombic phase. For x≧0.08, weak peaks of Co₃O₄ started appearingand gradually strengthened with increasing Co doping, suggesting asaturation limit of Co in SnO₂. It is noted that with increasing Coconcentration, the intensity of the cassiterite SnO₂ phase decreasedwhile the relative concentration of the orthorhombic phase graduallyincreased. Changes in the XRD peak intensity of the cassiterite phase(I_(c110)) and the orthorhombic phase fraction (x_(o)) of SnO₂ are shownin FIG. 9(a). The orthorhombic fraction x_(o) was calculated using themethod of standard additions,${x_{0} = \frac{K}{K + \left( {I_{c\quad 110}/I_{o\quad 111}} \right)}},$using K=2.69. Formation of the high-temperature orthorhombic SnO₂ phaseat ambient conditions has been observed in thin films and nanoscalepowders. Nucleation of the metastable orthorhombic phase has beenattributed to thin film strains and size-dependent internal pressuresdue to surface stresses in nanoparticles. Therefore, the increasingorthorhombic fraction of SnO₂ with Co concentration indicates that Codoping causes structural disorder and strain, and possible changes inthe particle size. The growth of the orthorhombic fraction is fast up to1% Co, above which a slower growth is observed, as shown in FIG. 9(a).This suggests that the intrinsic doping mechanisms active in the Coconcentration regimes above and below 1% may be different.

Average particle size L of the tetragonal SnO₂ phase was calculatedusing the width of the (110) peak and the Scherrer relation,$L = \frac{0.9\lambda}{B\quad\cos\quad\theta}$(where θ is the peak position, λ is the x-ray wavelength and B=(B_(m)²−B_(s) ²)^(1/2) was estimated using the measured peak width B_(m) andthe instrumental width B_(s)). These estimates showed that thecrystallite size decreased with increased Co doping, as shown in FIG.9(b). This, combined with the TEM images discuss below, indicates thatCo doping inhibits the growth of SnO₂ nanoparticles.

XRD peak positions showed significant changes with Co doping as shown inFIGS. 9(c) and 9(d). The tetragonal cassiterite SnO₂ peaks initiallyshifted to the higher 2θ angles as x increased to 0.01, as shown in FIG.9(d). But for x=0.03, there is a dramatic shift to the lower anglesfollowed by moderate changes in the peak positions at higher x. Thesechanges revealed interesting variations in the lattice parameters a andc with Co concentration as shown in FIG. 9(c). The SnO₂ latticeparameter a initially decreased due to Co doping for x≦0.01. Such arapid contraction of the lattice can be understood qualitativelyconsidering the sizes of the ions and their local coordinations.Substitution of 0.69 Å sized Sn⁴⁺ ions with 0.58 Å sized Co²⁺ ions isexpected to reduce the interatomic spacing significantly, justifying theinitial contraction of the lattice for x≦0.01. The observed rapidexpansion of the lattice for x=0.03 indicates a significantly differentdoping mechanism. At higher doping concentrations, the incorporation ofdopant ions in interstitial sites has been reported in some host systemscausing somewhat similar structural changes in the lattice parameters.

Interstitial incorporation of Co²⁺ ions might cause significant changesand disorder in the SnO₂ structure as well as many dramatic changes inthe properties of the material, which is discussed in the followingsections. The large difference in the charges and coordination numbersof Sn⁴⁺ and Co²⁺ ions will also contribute to the structural disorder inSnO₂ due to the removal of some oxygen ions that were attached to theoctahedrally coordinated Sn⁴⁺. For x≧0.03, the observed expansion alongthe a-direction and continued contraction along the c-direction shown inFIG. 9(c) will contribute to changes in the shape of the SnO₂ latticeand the nanoparticles. This was indeed observed in the TEM measurementsreported below. On increasing the Co concentration to 5%, the sphericalSnO₂ particles (observed in the undoped as well as 1% Co doped samples)appeared as nanorods with an aspect ratio as high as 3.

In FIG. 5(b), x-ray diffractograms (XRD) of Sn_(1-x)Co_(x)O₂ samplesprepared by annealing the dried precipitate at 600° C. for 6 hours areshown. All samples showed strong peaks due to the rutile-typecassiterite (tetragonal) phase of SnO₂ with relatively weaker peaks ofmetastable orthorhombic phase. Although cassiterite is the stable phaseunder ambient conditions, the orthorhombic phase also was observed inSnO₂ synthesized by various methods. The peak positions of the XRD linesof both phases did not show any measurable change, but the intensitiesof the peaks of the orthorhombic phase increased while that of therutile (cassiterite) SnO₂ phase decreased with increasing x, as shown inFIG. 5(b). As x increases, the XRD peaks appeared wider, indicatingpossible changes in the crystallite size and/or strain. No trace ofcobalt metal, oxides, or any binary tin cobalt phases, were observed inany of the samples doped with up to 5% Co, which is well above thedetection limit (˜1.5%) of the x-ray diffractometer used. At higher x,additional peaks of Co₃O₄ started appearing. The formation ofSn_(1-x)Co_(x)O₂ at high temperatures with x≦0.08 has been reportedusing a solid state reaction method, although at lower temperatures thesolubility was ≦2%. The wet chemical process used in this work isexpected to increase the solubility limit.

C. Optical Properties of the Cobalt-Doped Tin Dioxide

Room-temperature optical spectra in the ultraviolet and visible lightwavelength ranges were collected for the Sn_(1-x)Co_(x)O₂ samples usinga CARY 5000 spectrophotometer fitted with an integrating sphere diffusereflectance accessory. The spectrophotometer measures reflectancerelative to a background scatterer, which was powdered BaSO₄. Thesestudies indicated that samples with x≦1 do not have any Co₃O₄ phases.

Preliminary optical characterization of the pure and Co doped SnO₂powders were carried out by measuring the diffuse reflectance atroom-temperature. FIG. 10(a) shows a shift of the absorption edge tolonger wavelengths/lower energies and a decrease in the band gap of SnO₂for x≦0.01. Increasing the Co doping above this level reversed the trendin that the band gap increased gradually and a reduction in thereflectance was observed with Co percentage. The diffuse reflectance, R,of the sample is related to the Kubelka-Munk function F(R) by therelation F(R)=(1−R)²/2R, where R is the percentage reflectance. Thespectra used for the bandgap calculations are plotted in terms of F(R).The bandgap energies of the Sn_(1-x)Co_(x)O₂ powders were calculatedfrom their diffuse-reflectance spectra by plotting the square of theKubelka-Munk function, F(R)² vs. energy in electron volts. The linearpart of the curve was extrapolated to F(R)²=0 to get the direct bandgapenergy.

Diffuse reflectance measurements carried out on a pure nanoscale Co₃O₄reference sample prepared using an identical procedure (with x=1),showed prominent signatures at lower energies as shown in FIG. 10(b).Comparison of the optical spectra of Sn_(1-x)Co_(x)O₂ samples with thissuggests that samples with x<0.01 do not have any Co₃O₄ phase. However,evidence of its presence was observed in all samples with x≧0.03. Thisindicates that in samples with x=0.03 and 0.05, at least a fraction ofthe Co atoms precipitate as Co₃O₄, although XRD shows its formation onlyfor x≧0.08. Since the detection limit of the x-ray diffractometeremployed was ˜1.5% (determined from XRD measurements of the physicalmixtures of SnO₂ (x=0) and Co₃O₄ (x=1) nanoparticles prepared underidentical synthesis conditions), the fraction of Co forming Co₃O₄ in 3%and 5% Co doped SnO₂ should only be below this level. This explains whythe lattice parameters, band gap, particle size, shape, andorthorhombic/tetragonal fractions continue to change for x>0.01 althoughthe changes are relatively smaller in this range as illustrated in FIGS.9 and 10.

D. Raman Spectra of the Sn_(1-x)Co_(x)O₂ Samples

Raman spectra were collected for the Sn_(1-x)Co_(x)O₂ samples using aRenishaw S2000 Raman microscope. Samples were all probed using identicalinstrument conditions: 783 nm diode laser, 1200 line/mm grating, over aStokes Raman shift range of 50-1000 cm⁻¹. A line focus accessory wasalso employed, which permitted the collection of photon scatter datafrom an area ˜2 μm by 60 μm, rather than a discreet 1-2 μm diameterspot. Incident laser power was not measured; however, power at the laserhead was ˜28 mW, which would be expected to produce ˜2-4 mW at thesample. Sample preparation consisted of loosely packing the powder intoa stainless steel die accessory, which was then mounted on themicroscope stage for probing.

FIG. 11(a) shows the Raman spectra of Sn_(1-x)Co_(x)O₂ samples as afunction of Co concentration. The pure SnO₂ spectrum shown in FIGS.11(a) and 11(b) shows the classic cassiterite SnO₂ vibrations at 476cm⁻¹, 630 cm⁻¹, and 776 cm⁻¹. Addition of 0.5 or 1.0 molar percent Coresults in the appearance of two new peaks at 300 cm⁻¹ and 692 cm⁻¹.However, no significant change in the SnO₂ peak positions or widths wereobserved for these doping concentrations. These new Raman peaks at 300cm⁻¹ and 692 cm⁻¹ may be due to the vibrational modes activated by localstructural changes resulting from the substitution of Co²⁺ ions at theSn⁴⁺ sites. Further addition of Co to 3 or 5 molar percent results inthe appearance of peaks at 196 cm⁻¹, 480 cm⁻¹, 520 cm⁻¹, 617 cm⁻¹ and688 cm⁻¹. These five peaks observed in the samples with x≧0.03 matchwell with published Raman data of Co₃O₄. The appearance of Co₃O₄vibrational modes is in good agreement with the result from the opticaldata that predicts at least a fraction of the doped Co precipitates outas Co₃O₄ for x≧0.03

For Sn_(1-x)Co_(x)O₂ samples with x≧0.03, there is an apparentdisappearance of the SnO₂ peaks. This disappearance is most obvious forthe 630 cm⁻¹ SnO₂ peak as illustrated in FIG. 11(c) and would suggest aloss of the SnO₂ phase. However, the 776 cm⁻¹ SnO₂ peak is still visiblein the 3 and 5% Co doped samples. At best, its intensity has diminishedby a factor of 2. The extensive peak broadening and the subsequentdisappearance of the 630 cm⁻¹ peak and the loss in intensity of the 776cm⁻¹ peak are indicative of significant structural modifications anddisorder of the SnO₂ lattice for x>0.03. Based on the drastic changesobserved in the lattice parameter a, particle size and particle shape,this disappearance of the Raman peak also may be due to the interstitialincorporation of Co²⁺ ions and the subsequent structural changes.

As discussed above, at 0.5 and 1 molar percent of Co, a small Raman peakis present at 692 cm⁻¹. However, when the Co molar percentage isincreased to 3, an intense Raman peak at 688 cm⁻¹ appears. The width ofthe 688 cm⁻¹ peak precludes determination if the 692 cm⁻¹ is stillpresent. Two obvious conclusions are possible: i) the 692 cm⁻¹ peakrepresents a very small amount of Co₃O₄, and ii) the 692 cm⁻¹ moderepresents a vibrational mode of Sn_(1-x)Co_(x)O₂. FIG. 11(b) comparesthe Raman spectra of undoped SnO₂ and 1% Co doped SnO₂ (both prepared at600° C. through identical procedures) and the Raman spectrum of aphysical mixture of SnO₂ and Co₃O₄ with 1 molar percent Co (bothprepared separately at 600° C.). For the physical mixture, no annealingwas performed following the mixing. The Raman spectrum of the physicalmixture is clearly a superposition of a SnO₂ spectrum and a Co₃O₄spectrum. It is also clear that the intensity of the SnO₂ peaks in theundoped SnO₂, the 1% Co doped SnO₂, and the physical mixture samples,are essentially the same. This indicates that the same percentage ofSnO₂ in the cassiterite form is present in all three samples. Theintensity of the 692 cm⁻¹ peak in the 1% doped sample is roughly 1/10that of the 688 cm⁻¹ peak in the physical mixture. This disparity inintensity suggests that even if we assign the 692 cm⁻¹ peak to traces ofCo₃O₄, then there is no more than 0.1% Co₃O₄ in the 1% Co doped sample.

E. Shape and Size of the Particles in the Compositions

The shape and size of the particles in the compositions were determinedusing transmission electron microscopy. This showed that theSn_(1-x)Fe_(x)O₂ particles were all elongated with their average aspectrations changing from 1.25 and 70 nm long for x=0.01 to 1.7 and 25 nmlong for x=0.05. It is estimated that for ferromagnetic Sn_(1-x)FeO₂ inthis preferred embodiment, 95% of the particles are shorter than 100 nmlong. It is estimated that at least 95% of the Sn_(1-x)Co_(x)O₂particles are less than 50 nm long.

High-resolution transmission electron microscopy (TEM) analysis wascarried out on a JEOL JEM 2010 microscope with a specifiedpoint-to-point resolution of 0.194 nm. The operating voltage of themicroscope was 200 kV. All images were digitally recorded with a slowscan CCD camera (image size 1024×1024 pixels), and image processing wascarried out using the Digital Micrograph software from Gatan (Pleasant,California). Energy dispersive x-ray spectroscopy (EDX) was carried outusing the Oxford Link system attached to the TEM.

The transmission electron microscopy measurements showed significantchanges in the shape and size of the Sn_(1-x)Fe_(x)O₂ particlesdepending on the level of Fe-doping and the preparation temperature.Sn_(1-x)Fe_(x)O₂ particles prepared at 600° C. are shown in FIGS. 12(a)and 12(b). These particles were all elongated with their aspect ratiosand average length L changing from ˜1.25 and 70 nm, for x=0.01, to 1.7and 25 nm for x=0.05. Sn_(0.95)Fe_(0.05)O₂ particles annealed at 600° C.and 900° C. are shown in FIGS. 13(a) and 13(b). It is estimated that forferromagnetic Sn_(1-x)Fe_(x)O₂ in this preferred embodiment, 95% of theparticles are shorter than 100 nm long. These crystallite sizes matchvery well with similar estimates obtained from the XRD data, shown inFIG. 7(c). The energy dispersive x-ray spectroscopy measurements carriedout on 1% and 5% Fe-doped SnO₂ samples showed Fe concentrations inreasonable agreement with the estimates obtained from PIXE studies. TheTEM images of the Sn_(1-x)Fe_(x)O samples showed the presence of largemicron-sized particles with a different shape, as shown in FIG. 12(c).

TEM images also revealed significant differences in the shape of the600° C. prepared Sn_(1-x)CO_(x)O₂ particles doped with differentpercentages of Co. The Sn_(0.99)Co_(0.01)O₂ nanoparticles shown in FIG.5(c) were equiaxed (nearly spherical in shape) with an average diameterof 37 nm, which is about half the average size of the pure SnO₂particles (˜70 nm) prepared under similar conditions. In contrast, theSn_(0.95)Co_(0.05)O₂ particles shown in FIG. 5(d) were elongated intonanorods with an average aspect ratio of about 3. TheSn_(0.99)Cu_(0.01)O₂ particles had a higher average volume of 25000 nm³in comparison to the Sn_(0.95)Co_(0.05)O₂ particles (average volume˜10000 nm³). It is estimated that at least 95% of the particles are lessthan 50 nm long. Electron diffraction patterns taken from an aggregateof particles belonging to each of these samples, shown in the insets ofFIGS. 5(c) and 5(d), revealed characteristic ring patterns that confirmthe structure and phase purity measured by XRD. The differences in theparticle sizes and shapes observed in the TEM data as well as thesystematic changes in the linewidth and intensity of the XRD peaks withincreasing Co % seem to be the result of Co incorporation into the SnO₂lattice. The Co incorporation will most likely result in structuralrearrangements to take care of the ionic size differences and chargeneutrality. However, more detailed investigations are required to fullyunderstand these structural changes.

F. Mossbauer spectra of Sn_(1-x)Fe_(x)O₂

Mossbauer spectroscopy measurements showed that Sn_(0.95)Fe_(0.05)O₂exhibited ferromagnetically ordered Fe³⁺ spins when prepared at 350° C.,but that these ferromagnetically ordered Fe³⁺ spins were converted to aparamagnetic spin system as the preparation temperature increased to600° C. For these measurements, randomly oriented absorbers wereprepared by mixing approximately 30 mg of sample with petroleum jelly ina 0.375 inch thick and 0.5 inch internal diameter Cu holder sealed atone end with clear tape. The holder was entirely filled with the samplemixture and sealed at the other end with tape. Spectra were collectedusing a 50 mCi (initial strength) ⁵⁷Co/Rh source. The velocitytransducer MVT-1000 (WissEL) was operated in constant acceleration mode(23 Hz, ∓12 mm/s). An Ar—Kr proportional counter was used to detect theradiation transmitted through the holder, and the counts stored in amultichannel scalar as a function of energy (transducer velocity) usinga 1024 channel analyzer. Data were folded to 512 channels to give a flatbackground and a zero-velocity position corresponding to the centershift (CS or δ) of a metallic iron foil at room temperature. Calibrationspectra were obtained with a 20 μm thick α-Fe(m) foil (Amersham,England) placed in exactly the same position as the samples to minimizeany errors due to changes in geometry. Sample thickness corrections werenot carried out. The data were modeled with RECOIL software (Universityof Ottawa, Canada) using a Voigt-based spectral fitting routine.

Three selected samples, Sn_(0.95)Fe_(0.05)O prepared at 200° C.,Sn_(0.95)Fe_(0.05)O₂ prepared at 35° C., and Sn_(0.95)Fe_(0.05)O₂prepared at 600° C., were investigated using Mossbauer spectroscopy, andtheir spectra are shown in FIG. 14. The Sn_(0.95)Fe_(0.05)O sampleprepared at 200° C. showed a well-defined doublet in FIG. 14(a),suggesting that the incorporated Fe is paramagnetic and in the3+oxidation state in an octahedral environment. No evidence of any Fe²⁺ions was detected in this sample.

Experimental and fit-derived RT Mossbauer spectra of theSn_(0.95)Fe_(0.05)O₂ sample prepared at 350° C. are shown in FIG. 14(b).The spectrum displayed a well-defined sextet (magnetic componentspanning 24% of the spectral area) and a central doublet (paramagneticcomponent spanning 76% of the spectral area). The fit-derived parametersof the Fe sextet are center shift δ=0.38 mm/s, quadrupole shiftparameter ε=−0.1 mm/s, and hyperfine magnetic field B_(hf)=50.3 T,whereas the parameters of the doublet are δ=0.37 mm/s and quadrupoleshift ε=−0.1 mm/s.

FIG. 14(b) shows that the sextet feature (magnetic component of thesample) of the Sn_(0.95)Fe_(0.05)O₂ prepared at 350° C. is not due tocrystalline bulk Fe oxides such as magnetite, hematite, goethite, ormaghemite. Magnetite, a mixed oxide of Fe²⁺ and Fe³⁺, displays twowell-defined sextets in its RT Mossbauer spectrum because of thepresence of Fe in both tetrahedral (Fe³⁺) and octahedral sites (Fe²⁺ andFe³⁺ at a 1:1 ratio displays a sextet peak due to Fe^(2.5+) because ofVerwey transition): A and B sites of the inverse spinel structure.Hematite, on the other hand, displays a well-defined sextet withB_(hf)=51.8 T, δ=0.37 mm/s, and ε=−0.20 mm/s, which is not shown in FIG.14(b). Goethite displays a well-defined sextet with B_(hf)=38 T, δ=0.37mm/s, and ε=−0.26 mm/s, which is also not shown in FIG. 14(b). Thesextet feature shown in FIG. 14(b) is also unlikely to be due tomaghemite, which has B_(hf)=50 T, δ=0.23 to 0.35 mm/s, and ε=−0.02 mm/s.The experimental conditions employed to synthesize the binary oxides inthis research also implied nonformation of maghemite.

The derived Mossbauer parameters of the central doublet, which is due tocontribution from paramagnetic Fe site(s) to the sample, do not favorthe formation of small particle magnetite and goethite. Small-particleFe oxides such as magnetite (<10 nm), goethite (<15 nm), and hematite(<8 nm) display a doublet at room temperature (well below their magneticordering temperature) due to superparamagnetism. The parameters of thedoublet in FIG. 14(b), however, are inconsistent with superparamagneticiron oxides. The derived Mossbauer parameters of magnetite and hematiteare δ=0.22 mm/s, and ε=0 0.6 mm/s, and δ=0.35 mm/s, and ε=0.49 mm/s,respectively, while the quadrupole splitting of goethite is aroundε=0.48 mm/s. Moreover, any such iron oxides, if present in thesuperparamagnetic nanoparticle form, cannot produce the hysteresis loopswith a finite coercivity (˜60 Oe) observed in the magnetic studies. Thisobservation, along with the nonformation of “large” particle magnetite,hematite, and goethite in the sample, implies an absence of conditionsthat would dictate their formation. Therefore, it is believed that ironoxide does not exist in the preferred embodiment of iron-doped tindioxide.

Thus, the Mossbauer data shown in FIG. 14(b), showing the Mossbauerspectra of room-temperature Sn_(0.95)Fe_(0.05)O₂ prepared at 350° C.,suggest that the sextet results from magnetically ordered Fe³⁺, whichconstitutes 24% of the incorporated Fe ions. The Mossbauer spectra ofroom-temperature Sn_(0.95)Fe_(0.05)O₂ prepared by annealing theprecipitate at 600° C. shown in FIG. 14(c), on the other hand, showsmainly a doublet structure (corresponding to paramagnetic Fe³⁺) withvery weak traces of the sextet lines. This clearly suggests that theferromagnetically ordered Fe³⁺ spins are converted to a paramagneticspin system as the preparation temperature increased from 350 to 600° C.

E. The Distribution of the Dopant throughout the Crystallite

1. Confirmation by x-ray diffraction studies of Sn_(1-x)Fe_(x)O₂

X-ray diffraction (XRD) studies utilizing the Debye-Scherrer techniquehave also shown that the Fe is evenly distributed throughout the SnO₂,meaning that there are no phases or clusters of the Fe. The pure ironoxide samples (prepared under identical synthesis conditions, but withno SnCl₂) showed maghemite [γ-Fe₂O₃, FIG. 6(a) and hematite [α-Fe₂O₃,FIG. 6(b) phases, with average particle sizes of 22 and 53 nm at 200 and600° C. preparations, respectively. No trace of iron metal, oxides, orany binary tin-iron phases were observed in any of the doped sampleswith x≦0.10. This is consistent with the reported solubility limit of upto 10% Fe in SnO₂ (this, along with the fact that the x-raydiffractometer employed can detect only phases that are ≧1.5%, is whyexperimental data has been reported only for samples with x≦0.05). Thelack of any observable phases shows that the Fe is evenly distributedthroughout the SnO₂.

2. Confirmation by X-Ray Photoelectron Spectroscopy Studies

X-ray photoelectron spectroscopy studies (XPS) showed thatSn_(1-x)Fe_(x)O₂ and Sn_(1-x)Co_(x)O₂ prepared according to the methodsdisclosed herein produce a uniform distribution of the dopant in theentire crystallite. XPS measurements for both Fe-doped and Co-doped SnO₂were performed on powder samples using a Physical Electronics Quantum2000 Scanning ESCA Microprobe. The system used a focused monochromaticA1K_(α) x-ray (1486.7 eV) source and a spherical section analyzer. Theinstrument had a 16 element multichannel detector. The x-ray beam usedwas a 105 W, 100 μm diameter beam that was rastered over a 1.4 mm×0.2 mmrectangle on the sample. The powder samples were mounted using a smallamount of double-coated carbon conductive tape. The x-ray beam wasincident normal to the sample and the photoelectron detector was at 45°off-normal. Data were collected using a pass energy of 46.95 eV. For theAg 3d_(5/2) line, these conditions produce full width at half-maximum ofbetter than 0.98 eV. Although the binding energy (BE) scale wascalibrated using the Cu 2p_(3/2) feature at 932.62∓0.05 eV and Au 4ffeature at 83.96∓0.05 eV for known standards, both the Fe-doped andCo-doped SnO₂ surfaces experienced variable degrees of charging.Low-energy electrons at ˜1 eV, 21 μA, and low-energy Ar⁺ ions were usedto minimize this charging. The BE positions were referenced using the486.7 eV position for the Sn 3d_(5/2) feature for the Sn_(1-x)Fe_(x)O₂samples and for the Sn_(1-x)Co_(x)O₂ samples, and the 486.9 eV positionfor the Sn_(1-x)Fe_(x)O samples. XPS spectra were also collected afterAr⁺ ion sputtering using a 4 kV Ar⁺ ion beam rastered over a 4 mm×4 mmsample area. The sputter rates were calibrated using a SiO₂ standardwith known thickness.

The Fe 3p_(1/2) XPS spectral region of Sn_(1-x)Fe_(x)O₂ (prepared byannealing the precipitate at 600° C.) samples with x=0.01 and 0.05 areshown in FIG. 15. Similar data obtained from hematite and maghemitephases of pure Fe₂O₃ prepared under identical conditions are also shownfor comparison. The XPS peaks are not clearly visible in the 1% Fe-dopedsamples; however, clear peaks were observed in Sn_(0.95)Fe_(0.05)O₂.This discrepancy may be related to the limited Fe detection ability ofthe XPS system. The XPS peaks of both hematite and maghemite phasesoccur at ˜55.7 eV, which matches well with literature reports. However,the core level peak of Fe in Sn_(1-x)Fe_(x)O₂, as shown in FIG. 15, andSn_(1-x)Fe_(x)O, as shown in FIG. 16, 200° C. data, showed a slightshift to higher binding energies (˜56.5 eV) compared to the Fe oxides,indicating the difference in the atomic environment surrounding theincorporated Fe ions. Fe 3p_(1/2) has a reported binding energy of 53.9eV for the magnetite (Fe₃O₄) form of iron oxide. Thus, the XPS dataclearly suggest that the Fe peaks observed from the Sn_(1-x)Fe_(x)O₂ andSn_(1-x)Fe_(x)O samples are not arising from any maghemite, hematite, ormagnetite inclusions in the samples. The relative peak positions of theSn and O peaks in the samples did not show any measurable change with Feconcentration, suggesting that their chemical environments did notchange significantly. Atomic percentages of Sn, Fe, and O calculatedusing the Sn 3d_(5/2), Fe 3p_(1/2), and O 1s peaks are given in Table 1.

FIG. 16 shows the XPS data obtained from the Sn_(1-x)Fe_(x)O andSn_(1-x)Fe_(x)O₂ samples prepared by annealing the same reactionprecipitate at 200, 350, 450, 600, 750, and 900° C. In all thesesamples, the core level Fe peak was observed at ˜56.5 eV and nomeasurable shifts towards the binding energies expected for magnetite(53.9 eV), hematite (55.7 eV), and maghemite (55.7 eV) were observedwhen the preparation temperature varied in the 350 to 900° C. range.Although the Fe doping concentration was 5%, the Fe XPS peakssystematically intensified with increasing preparation temperature.Atomic percentages of Sn, Fe, and O, calculated using the XPS peaks as afunction of preparation temperature, are given in Table I.Notwithstanding the difference between the atomic concentrationsobtained from PIXE and XPS, the XPS estimates from Sn_(0.95)Fe_(0.05)O₂showed a systematic increase in the Fe concentration from 0.7% to 6.2%as the annealing temperature increased from 350 to 900° C. As mentionedabove, the lower Fe estimates from the XPS data may be due to therelatively lower detectability of Fe using XPS. In the PIXE measurementsdiscussed above, the Fe concentration of the Sn_(1-x)Fe_(x)O₂ samplesprepared in the entire temperature range was always between 4% and4.88%, and no systematic variation with preparation temperature wasobserved. Compared to PIXE, which is responsive to the entire bulk ofthe sample, XPS is a surface-sensitive technique. Therefore, theincreasing differences between the Fe concentrations obtained from thesetwo techniques clearly suggest a gradual and systematic surfacediffusion of the doped Fe ions with increasing preparation temperature.This suggests that the Sn_(1-x)Fe_(x)O₂ samples prepared at lowertemperature produce a more uniform distribution of Fe in the entirecrystallite. On the other hand, samples prepared at higher temperaturesshowed significant diffusion of the incorporated Fe ions towards theparticle surface as preparation temperature increases to 900° C.

To further confirm the Fe surface diffusion possibility, XPS spectrawere collected from the 900° C. prepared Sn_(1-x)Fe_(x)O₂ sampleemploying Ar⁺ ion sputtering to remove surface layers from the powdersamples mounted on carbon conductive tape, as shown in FIG. 17. Thesemeasurements showed a gradual decrease in the Fe concentration from6.23% in the as-prepared sample to 3.43% when a 20 nm surface layer isremoved by Ar⁺ ion sputtering (see Table I). This fully supports theabove conclusion that the higher XPS estimates of Fe concentrationobtained from Sn_(1-x)Fe_(x)O₂ samples prepared at higher temperatures(≧600° C.) are indeed due to Fe surface diffusion.

The Co 2p_(3/2) and Co 2p_(1/2) XPS spectral region of theSn_(1-x)Co_(x)O₂ samples are shown in FIG. 18. Comparing the bindingenergies of the Co primary and satellite XPS peaks with that observedfor Co(0) in Co metal, Co²⁺ in CoO and Co³⁺ in δ-Co₂O₃, the electronicstate of Co in Sn_(1-x)Co_(x)O₂ samples is found to be Co²⁺ and that itis not bonded to oxygen as CoO or Co₃O₄. It also rules out any metallicCo clusters in the samples, a result well expected for chemicallysynthesized samples prepared and processed in air. These results agreewell with the 2+ oxidation state of Co with S=3/2 determined frommagnetization measurements of paramagnetic samples of Sn_(1-x)Co_(x)O₂.The >1 eV shift of the Co 2p_(3/2) peak in the Sn_(1-x)Co_(x)O₂ samplescompared to that observed from the Co₃O₄ reference sample suggests thatCo is indeed incorporated in the SnO₂ lattice and not forming anysignificant amount of Co oxides. However, no significant change in theCo binding energy is observed with increasing Co doping concentration.Careful analysis of the peak positions of the Sn 3d_(5/2) (486.7 eV) andO 1 s (530.65 eV) peaks also did not show any noticeable change in thebinding energy with increasing Co concentration.

Atomic percentages of Sn, Co, and O calculated using the Sn 3d_(5/2)(486.7 eV), O 1 s (530.65 eV), and Co 2p_(3/2) (781.4 eV) peaks areshown in FIG. 19. In these plots, three well defined regions arepresent. A rapid increase in the atomic percentage of Co and the Co/Snratio, and a decrease in the Sn atomic percentage for x≦0.01, indicatesubstitutional incorporation of Co at Sn sites. In the 0.01≦x≦0.05range, a relatively slower variation is observed suggesting moreinterstitial incorporation and/or Co₃O₄ precipitation in agreement withresults from XRD, Raman, and optical studies, discussed above. In thecase of substitutional Co incorporation, Co ions remove Sn ions from theSnO₂ structure. However, in the case of interstitial incorporation orCo₃O₄ precipitation, there is no Sn removal. For x>0.05, changes in theCo/Sn atomic percentage ratio are minimal, indicating lack of further Coincorporation into the SnO₂ lattice. Estimation of the oxygen content inthe samples using the O 1 s (530.65 eV) peak indicated almoststoichiometric Sn/O ratio for the undoped sample. Co doping decreasesthe oxygen content of the sample, as shown in FIG. 19(a). Removal ofoxygen atoms from the SnO₂ lattice is well expected if Co²⁺ replacesSn⁴⁺ ions due to charge neutrality requirements. It has been argued thatsubstitutional Co²⁺ and oxygen vacancies in excess of those necessaryfor charge neutrality are essential to produce ferromagnetism inTi_(1-x)Co_(x)O₂. A rapid loss in oxygen content, indicated byrelatively larger changes in the oxygen atomic percentage, for sampleswith x≦0.01 (FIG. 19(a)), may be crucial for the ferromagnetism observedonly in this narrow Co percentage range of ≦1%.

3. Shown by Absence of Iron Oxide Phases in the Sn_(1-x)Fe_(x)O₂

The even distribution of Fe throughout the ferromagneticSn_(1-x)Co_(x)O₂ powder has been shown by the absence of iron oxidephases in the samples. The origin of ferromagnetism in dilute magneticsemiconductor oxides has been extensively studied recently because ofthe possible presence of weaker secondary phases. This is particularlyimportant when the ferromagnetic component is weak. The fact that thesol-gel preparation of the samples and their subsequent drying andannealing processes were all conducted in air intrinsically eliminatesthe possibility of forming metallic Fe particles.

The possibility was investigated that the ferromagnetism observed inSn_(1-x)Co_(x)O₂, when prepared in the 350 to 600° C. range, may be dueto weak traces of maghemite or magnetite phases of iron oxide formed inthe sample. The pure iron oxide samples prepared under identicalsynthesis conditions showed the formation of pure maghemite whenprepared at 200° C. and pure hematite at 600° C. However, noferromagnetism was observed in the Sn_(1-x)Fe_(x)O sample prepared byannealing the precipitate at 200° C., which rules out the presence ofany maghemite phase undetected in the XRD data. Therefore, it isunlikely that this phase will appear when the Sn_(1-x)Fe_(x)O₂ sample isprepared by annealing the same precipitate in the 350 to 600° C. range.Investigation of the phase transition of pure iron oxide samplesprepared under identical conditions showed that the maghemite phaseconverted to the hematite phase when annealed at temperatures above 350°C. Thus, it is very unlikely that the maghemite phase of iron oxide ispresent in the Sn_(1-x)Co_(x)O₂ samples prepared by annealing attemperatures ≧350° C.

The M versus H data shown in FIG. 20(c), obtained from measuringSn_(1-x)Co_(x)O₂ at 5 K, and M versus T data shown in FIG. 21, obtainedfrom measuring pure iron oxide samples prepared at 200° C. and 600° C.with an applied magnetic field H=500 Oe, ruled out the presence ofhematite due to the absence of spin-flop and the strong Morintransitions. The hematite phase is thermally the most stable phase andit undergoes a thermal reduction to the magnetite (Fe₃O₄) phase onlyabove 1200° C. Thus, thermodynamically, the possible formation of themagnetite phase can also be ruled out. Even if the magnetite phase wereformed, the observed disappearance of ferromagnetism when prepared attemperatures above 600° C., as shown in FIG. 22, would be difficult tounderstand.

Further, careful analysis of the samples using XRD, TEM, and selectedarea diffractions experiments has ruled out the presence of any ironoxide phases in the Sn_(1-x)Co_(x)O₂ samples. Finally, the Mossbauerdata, XPS spectra, and hysteresis loop parameters obtained from theSn₁Fe_(x)O₂ samples clearly ruled out the presence of any bulk ornanoscale magnetite, hematite, maghemite, or goethite phases of ironoxide in the samples.

Room-temperature ferromagnetism observed in a Mn-doped ZnO system hasbeen shown to result from a metastable Mn_(2-x)ZnXO_(3-d) type phaseformed by the diffusion of Zn into Mn oxides. In these studies, peaksdue to pure and/or doped manganese oxides were clearly observed in theXRD measurements (plotted on a log scale). In the present work, althoughthe saturation magnetization increases by about four times as Feconcentration increases to 5%, no indication of pure or doped ironoxides or other impurities is observed in the XRD measurements (shown onlog scale), as illustrated in FIG. 23. In the selected area electrondiffraction, XPS, or Mossbauer spectroscopy studies as well, no evidencefor the presence of any such mixed phases in Sn_(1-x)Fe_(x)O₂ in theentire ranges of Fe concentration and preparation temperatures wereobserved. The sol-gel-based chemical synthesis employed to manufacturethe Sn_(1-x)Fe_(x)O₂ is well known to provide a uniform distribution ofthe dopant in the host system at lower temperatures, as compared to thesolid state reaction used in the preparation of Mn-doped ZnO.

4. Incorporation of Fe into the SnO₂ and SnO Lattices

The systematic changes in the lattice parameters, particle size, andshape observed in XRD and TEM studies strongly support the progressiveincorporation of Fe into the SnO₂ and SnO lattices with increasing x.The one-to-one match in the relative changes in the saturationmagnetization M_(s) and lattice volume V, shown in FIG. 24, observed inthe Sn_(1-x)Fe_(x)O₂ samples, is very strong evidence against anyimpurity being the origin of the observed ferromagnetism. The one-to-onematch also suggests a strong structure-magnetic property relationship inthese samples. The striking agreement between the estimated magneticallyordered Fe³⁺ spins (˜24%) in the powder samples of Sn_(0.95)Fe_(0.05)O₂and a similar estimate of ferromagnetic Fe³⁺ spins (˜23%) inpulsed-laser-deposited thin films of Sn_(0.95)Fe_(0.05)O₂ furthersupports the conclusion that the observed ferromagnetism inSn_(1-x)Fe_(x)O₂ is not due to impurity iron oxide phases formed underthe different preparation conditions employed.

The conclusion that the Fe is incorporated into the SnO₂ and SnOlattices is also supported by the role of the host system. It is wellknown that the p-type semiconducting behavior SnO results from an excessof oxygen, whereas the existence of oxygen vacancies in SnO₂ make it anexcellent n-type semiconductor. The XPS data obtained for 1% and 5%Fe-doped SnO showed identical oxygen atomic percentages (see Table I),whereas the oxygen concentration decreased in Sn_(1-x)Fe_(x)O₂ with Feconcentration. The Sn-O distance of 2.057 Å in SnO₂ is lower than the2.223 Å in SnO, and this might influence the overlap of the electronorbitals. Thus, in Sn_(1-x)Fe_(x)O, Fe doping might favor the formationof antiferromagnetic Fe³⁺—O²⁻Fe³⁺ groups, whereas Sn_(1-x)Fe_(x)O₂ willhave a large number of ferromagnetic Fe³⁺-[oxygen vacancies]-Fe³⁺ groupsbecause of the oxygen vacancies. This might explain the observedantiferromagnetic interaction in Sn_(1-x)Fe_(x)O and ferromagnetism inSn_(1-x)Fe_(x)O₂.

The Sn_(1-x)Fe_(x)O₂ composition showed a strong structure-magneticproperty relationship, as shown in FIG. 24(a), where the increase in thesaturation magnetization with Fe concentration matches with the increasein the lattice concentration. Sn_(1-x)Fe_(x)O, on the other hand, showedan expansion of the lattice with increasing Fe concentration, and hereno ferromagnetism is observed. Changes in the internal or externallattice volume/pressure have been reported to produce ferromagnetism initinerant electron metamagnets. Thus, more investigation is required tounderstand the exact role of structural changes and internal pressuredifferences in the observed ferromagnetism/paramagnetism ofSn_(1-x)Fe_(x)O₂/Sn_(1-x)Fe_(x)O.

F. Magnetic Properties of the Compositions

The Sn_(1-x)Fe_(x)O₂ showed ferromagnetic behavior with a Curietemperature of up to 850 K, well above room-temperature, for the 1%Fe-doped sample. All of the Sn_(1-x)Fe_(x)O₂ samples show well-definedhysteresis loops at 300 K, room-temperature, with remanence M_(r) andsaturation magnetization M_(s) increasing gradually with the level ofFe-doping. The ferromagnetic property is stronger when prepared at lowerannealing temperatures, and it gradually declines with increasingpreparation temperature and eventually disappears completely forpreparation temperatures greater than 600° C. In the preferredembodiment, the Sn_(1-x)Fe_(x)O₂ powder is free of any hematiteparticles.

Magnetic measurements for both Sn_(1-x)Fe_(x)O₂ and Sn_(1-x)CO_(x)O₂were carried out as a function of temperature (4 to 350 K) and magneticfield (0 to ˜65 kOe) using a commercial magnetometer (Quantum Design,PPMS) equipped with a superconducting magnet. Measurements were carriedout on tightly packed powder samples placed in a clear plastic drinkingstraw. The data reported were corrected for the background signal fromthe sample holder (clear plastic drinking straw) with diamagneticsusceptibility χ=−4.1×10 ⁻⁸ emu/Oe.

1. Iron Concentration Dependence

a. Sn_(1-x)FeO₂

The room-temperature M versus H data of Sn_(1-x)Fe_(x)O₂, shown in FIG.20(a), show a linear component superimposed on a saturatingferromagnetic-like magnetization. If this linear component χ_(p) issubtracted, the M−χ_(p) data show saturation of M expected for aferromagnetic phase, as shown in FIG. 20(b). FIG. 20(c) shows that at 5K, the ferromagnetic component is overwhelmed by a paramagnetic-likecomponent. Variations of the saturation magnetization Ms and χ_(p)obtained from the M versus H data as a function of the Fe-doping areshown in FIGS. 24(a) and 24(b). These data fit reasonably well with themodified Brillouin function, assuming that J=5/2. This indicates that afraction of the doped Fe is not participating in the ordered magneticstate, in excellent agreement with the Mossbauer results previouslyshown in FIG. 14(b). The exact nature of this component becomes moreevident from the M versus T data shown in FIG. 25. This showed aparamagnetic variation described by the modified Curie-Weiss law similarto their Sn_(1-x)Fe_(x)O counterparts as shown in FIG. 26, but offset byan amount χ₀; this offset is most likely due to the ferromagneticcomponent. This also confirms that the linear component χ_(p) observedin the room-temperature M versus H data shown in FIG. 20(a) is also dueto this paramagnetic contribution present in the sample. T₀ and θ data,shown in FIG. 24(c), obtained from the M versus H and M versus T datarespectively, indicate that the interaction between the disordered(paramagnetic-like) Fe³⁺ spins present in Sn_(1-x)Fe_(x)O₂ isantiferromagnetic (AF) in nature.

Measurements of the sample magnetization M as a function of magneticfield H and temperature T were carried out using a commercialmagnetometer (Quantum Design, PPMS) equipped with a superconductingmagnet. The data reported were corrected for the background signal fromthe sample holder. In the inset of FIG. 27, M vs. H plots ofSn_(1-x)Co_(x)O₂ samples (with x=0.01 and 0.03) prepared at 600° C. andmeasured at 5K are shown along with their theoretical estimates obtainedusing the Brillouin-function-based form for a paramagnetic system, givenbyM=M ₀{[(2J+1)/2J]coth[(2J+1)y/(2J)]−(½J)coth(y/2J)}Where y=(gμ_(B)JH)/(kT), M₀ is the saturation magnetization, g is thespectroscopic splitting factor (g=2.0023 for free electrons), μ_(B) isthe Bohr magneton and k is the Boltzmann constant. M vs. H data of theSn_(1-x)Co_(x)O₂ samples fit very well with their theoretical estimatesyielding a total angular momentum J=1.81±0.1. These values are in goodagreement with experimental magnitudes (˜4.8) reported for paramagneticCo²⁺ ions with spin S=3/2 [13].

Magnetic susceptibility χ=M/H of the samples measured as a function oftemperature at a constant H=500 Oe also showed the expected paramagneticbehavior. In FIG. 27, X vs. T data of the 1 and 3% Co doped samples areshown along with theoretical fits obtained using the modifiedCurie-Weiss lawX=X ₀ +[C/(T+θ)]Where X₀=1.5(0.2)×10⁻⁶ emu/g Oe represents weak non-paramagneticcontribution, Curie constant C=Nμ²/3k is a measure of the paramagneticion concentration (N=number of magnetic ions/g, μ=magnetic moment of theion) and θ is the Curie-Weiss temperature which represents the magneticexchange interactions between the spins. These fits yield θ=0.18 and1.55K, and C=0.63×10⁻⁴ and 1.6×10⁻⁴ emu K/g Oe for x=0.01 and 0.03respectively.

The pure hematite form of iron oxide, prepared at 600° C. following anidentical synthesis procedure but with no Sn precursor, showed a weakmagnetization, as shown in FIGS. 20(a), 20(c), and 21. The most strikingcharacteristics of bulk hematite include the sharp Morin transition near263 K in the M versus T data and a spin-flop (SF) transition atH_(SF)˜67.5 kOe in the M versus H data. Both of these transitions wereindeed present in our pure hematite as shown in FIGS. 20(a) and 21,albeit with reduced magnitudes which are presumably due to a smallerparticle size of ˜53 nm. These transitions were clearly absent in all ofthe Sn_(1-x)Fe_(x)O₂ samples, ruling out the presence of any hematiteparticles.

The Sn_(1-x)Fe_(x)O₂ samples showed well defined hysteresis loops at 300K, as shown in FIG. 28, with remanence M_(r) and saturationmagnetization M_(s) increasing gradually with the percentage ofFe-doping, as shown in FIGS. 24(a) and 24(b). The coercivity Hc was inthe range of ˜60 Oe, which is significantly different from the value ofH_(c)=1844 Oe obtained for the pure hematite sample prepared underidentical conditions. The existence of a significant coercivity inSn_(1-x)Fe_(x)O₂ samples clearly rules out their possible origin fromnanoscale superparamagnetic particles of iron oxides because whenmagnetic materials are prepared in nanoscale sizes, they demonstratesuperparamagnetic behavior characterized by hysteresis loops with zerocoercivity above their blocking temperatures. Absence of bulk ironoxides (or nonsuperparamagnetic particles) was confirmed from theMossbauer data discussed above. Finally, a Curie temperature T_(c)=850 Kwas obtained for the 1% Fe-doped SnO₂ sample by measuring M up to 1000K, as shown in FIG. 29. This Curie temperature is among the highestCurie temperatures reported for transition-metal-doped oxidesemiconductors.

Some of the samples with 0.5 and 1% Co doping annealed at 600° C. showeda ferromagnetic behavior. In the inset (a) of FIG. 30, a hysteresis loopmeasured at 300 K from a 1% Co doped SnO₂ sample prepared by annealingthe precipitate at 600° C. is shown. The magnetization saturates verywell for H>3 kOe with a saturation magnitude of 0.133 μ_(B)/Co ion.However, the data did not show any significant coercivity or remenance.

The systematic growth of both ferromagnetic and paramagneticcontributions in Sn_(1-x)Fe_(x)O₂ with increasing x, as shown in FIGS.24(a) and 24(b), suggests that the ferromagnetic component is notgrowing at the expense of the paramagnetic Fe³⁺ ions as Fe dopingincreases. Other researchers have proposed a ferromagnetic exchangemechanism involving oxygen vacancies, which form F-centers with trappedelectrons, for the observed ferromagnetism in Fe-doped SnO₂ thin films.Overlap of the F-center electron orbitals with the d-orbitals of theneighboring Fe³⁺ spins to form Fe³⁺-[oxygen vacancies]-Fe³⁺ groups iscrucial for the proposed ferromagnetic coupling. It has been argued thatdoped Fe³⁺ spins might also exist as isolated paramagnetic spin systemswherever the F-center mediated ferromagnetic coupling is not achieveddue to lack of Fe³⁺ neighbors and/or oxygen vacancies. In addition, anyFe³⁺—O² ⁻—Fe³⁺ superexchange interactions will be antiferromagnetic innature. As Fe doping concentration increases, both ferromagnetic andparamagnetic/antiferromagnetic components will increase leading to theobserved variations shown in Figures (FIGS. 15(a) and 15(b)). It isbelieved that Sn_(1-x)Fe_(x)O₂ will exhibit ferromagnetism for any valueof x up to the solubility limit of Fe in SnO₂, or 10%.

b. Sn_(1-x)Fe₁O

FIG. 31(a) shows the magnetization M of the Sn_(1-x)Fe_(x)O samplesmeasured at 5 K as a function of applied magnetic field H along withtheir theoretical estimates obtained using the modifiedBrillouin-function-based form for a paramagnetic system, given byM=M ₀{[(2J+1)/2J]coth[(2J+1)y/(2J)]−(½J)coth(y/2J)}

where y=gμ_(B)JH/k(T+T₀), M₀ is the saturation magnetization, g=2.0023is the spectroscopic splitting factor, μ_(B) is the Bohr magneton, and kis the Boltzman constant. Based on the Mossbauer data discussed above,this analysis was carried out assuming that J=5/2 (which is expected forFe³⁺). Here, T₀ is included as a measure of the magnetic interactionbetween the Fe spins, which prevents complete alignment of the spinseven at the highest magnetic fields employed. A larger T₀ indicatesstronger antiferromagnetic (AF) interactions between the disordered Fespins. Magnitudes of M₀ and T₀ obtained from this analysis are shown inTable II. M versus H plots of Sn_(1-x)Fe_(x)O samples measured at 300 Kshowed a linear variation owing to the paramagnetic behavior, as shownin FIG. 31(b). TABLE II Variations of magnetization M₀, interactiontemperature T₀, Curie constant C, and Curie-Weiss temperature θ ofSn_(1−x)Fe_(x)O as a function of x. Sn_(1−x)Fe_(x)O Fe MagnetizationInteraction Curie constant Curie-Weiss percentage M₀ (emu/g) temperatureT₀ (K) (10⁻⁴ emu K/g Oe) Temperature θ (K) 1 0.25 4.50 0.55 4.17 5 1.243.10 2.81 3.00

Magnetization M of the Sn_(1-x)Fe_(x)O samples measured as a function oftemperature T at a constant field H=500 Oe also showed the expectedparamagnetic behavior, as shown in FIG. 26, following the modifiedCurie-Weiss law χ=χ₀+C/(T+θ), where χ₀=4(3)×10⁻⁶ emu/g Oe representsweak nonparamagnetic contributions, Curie constant C=Nμ²/3k (N=number ofmagnetic ions/g, μ=magnetic moment of the ion), and θ is the Curie-Weisstemperature. These fits showed an increase in C (as well as M₀) with x,as shown in Table II, confirming the progressive doping of Fe ions. Thepositive values of θ indicate AF interactions between the Fe spins asobserved in other systems as well. Both θ and T₀ decrease with x, asshown in Table II, indicating that the AF interaction decreases withincreasing Fe doping. This may suggest that there are competing AF andferromagnetic interactions as x increases.

The pure iron oxide sample prepared under identical conditions asSn_(1-x)Fe_(x)O was strongly ferromagnetic, as shown in FIGS. 31(a) and31(b). M versus T data, shown in FIG. 21, of this sample indicated ablocking temperature T_(B)˜21 K, suggesting the presence of nanoscaleferromagnetic particles. These observations match very well with the XRDdata showing the formation of nanoscale maghemite. This also rules outthe presence of this phase in the Sn_(1-x)Fe_(x)O samples, which are allparamagnetic for x≦0.05.

2. Cobalt Concentration Dependence

Magnetic measurements carried out on pure SnO₂ nanoparticles showed theexpected diamagnetism with a negative magnetic susceptibility. Applicanthas shown that the Sn_(1-x)Co_(x)O₂ samples with x≦0.01 were allferromagnetic at room-temperature when prepared in the 350 to 600° C.temperature range. FIG. 32(a) shows the room temperature hysteresis loopmeasured from a Sn_(0.99)Co_(0.01)O₂ sample prepared at 600° C.illustrating a clear ferromagnetic behavior with a coercivity H_(c)=9Oe. In FIG. 32(b), the variations of the room-temperature coercivityH_(c) and remanence M_(r) of 350° C. prepared Sn_(1-x)Co_(x)O₂ as afunction of Co concentration are shown. The observed coercivities of˜630 Oe and remanences as high as 31% are among the highest reported fordilute magnetic semiconductors. The observed variation of the saturationmagnetization Ms with Co concentration measured from the Sn_(-x)Co_(x)O₂samples prepared at 350 and 600° C. were comparable as illustrated inFIG. 6 c. However, the coercivities of the Sn_(0.99)Cu_(0.01)O₂ samplesdecreased from 630 Oe to 9 Oe as the preparation temperature increasedfrom 350 to 600° C. (see FIGS. 32(a) and 32(b)). This most likelyindicates a change in the magnetocrystalline anisotropy due to reasonsthat are unclear at present. More importantly, irrespective of thepreparation temperature, all Sn_(1-x)CO_(x)O₂ samples prepared in the350 to 600° C. showed complete destruction of the ferromagnetism above1% Co doping and only a paramagnetic behavior was observed in thisrange, as shown in FIG. 32(b). This disappearance of ferromagnetismcannot be explained by assuming a 1% solubility limit for Co in SnO₂because—(i) the ferromagnetic component due to the soluble part of thedoped Co should not be destroyed or overwhelmed by the weak paramagneticcomponent of the segregated Co₃O₄ formed for x>0.01, (ii) theorthorhombic SnO₂ fraction in the samples, lattice parameters, band gapenergy, particle size, and shape of the Sn_(1-x)CO_(x)O₂ particles(shown in FIGS. 9 and 10) continued to change with x for x>0.01,indicating >1% Co solubility in SnO₂, (iii) the observed disappearanceof the Raman peaks (FIG. 11(c)) for x>0.01 is unlikely to happen ifadditional Co doping is not taking place, and (iv) no evidence of anychange in the oxidation state of Co or Sn is observed in the XPSmeasurements for x>0.01.

The appearance of ferromagnetism in Sn_(1-x)Co_(x)O₂ samples with x≦0.01and its complete absence at higher Co concentrations can bequalitatively understood by comparing the changes in the magnetic andstructural properties noticed in the XRD, Raman, and TEM studies of the600° C. prepared samples. As shown in the previous sections, for x<0.01the SnO₂ lattice contracts, resulting in the reduction of the distancebetween nearby Co²⁺ spins, and possibly triggering a ferromagneticcoupling. Substitution of Sn⁴⁺ ions (octahedrally coordinated with sixnearest oxygen neighbors) in SnO₂ with Co²⁺ ions will result in thecreation of oxygen vacancies and additional charge carriers. It is notclear if this ferromagnetic ordering is carrier mediated or via othermechanisms such as based on localized defects (F-centers). Increasingthe Co doping to ≧3% results in a rapid expansion of the SnO₂ latticeand significant structural disorder indicated by the rapid broadeningand disappearance of the Raman peaks, as shown in FIG. 11(b). Suchenormous structural changes might have destroyed the ferromagneticordering since the magnetic exchange interaction is extremely sensitiveto the distance between the interacting spins.

It may be noted that the ferromagnetic regime of Sn_(1-x)Co_(x)O₂ withx≦0.01 corresponds to the compositions for which the SnO₂ latticecontracts (see FIGS. 32(b) and 32(c)). This might suggest that theobserved ferromagnetism may be related to internal pressure changes.Changes in the internal or external lattice volume/pressure have beenreported to produce ferromagnetism in itinerant electron metamagnets.

3. Temperature Dependence in Sn_(1-x)Fe_(x)O₂

The ferromagnetic component of Sn_(1-x)Fe_(x)O₂ gradually declines andsubsequently disappears as the preparation temperature increases, asshown in FIGS. 14, 22, and 33. Annealing the reaction precipitate attemperatures between 350 and 900° C. produces the Sn_(1-x)Fe_(x)O₂phase. The M versus H data measured from the Sn_(1-x)Fe_(x)O₂ preparedby annealing the same reaction precipitate at different temperature,shown in FIG. 33, clearly shows the presence of a ferromagneticcomponent in samples annealed at 350, 450, and 600° C. Only a linearvariation indicating a purely paramagnetic behavior was observed in thesample prepared by annealing at 750 and 900° C. The saturationmagnetization M_(s), estimated after subtracting the linear paramagneticcomponent χ_(p), is plotted in FIG. 22. This clearly establishes thefact that the ferromagnetic component is stronger when prepared at lowerannealing temperatures and it gradually decreases with increasingpreparation temperature, eventually disappearing completely forpreparation temperatures >600° C., which is in excellent agreement withthe Mossbauer data discussed above. The remanence M_(r) obtained fromthe hysteresis loops, shown in FIG. 22, also decreases with preparationtemperature. This figure also shows that for Sn_(1-x)Fe_(x)O₂ preparedat or below 600° C., the surface concentration of Fe is less than 4%.

Based on the observed changes in the Fe XPS peak intensity shown in FIG.16 and the comparison of the concentrations estimated from the PIXE andXPS data listed in Table I, it was concluded that the Fe ions diffusetowards the particle surface as the preparation temperature increases.The lattice volume plotted in FIG. 8(b) shows a gradual contraction ofthe lattice as preparation temperature increases, presumably due to theoutward diffusion and rearrangement of the doped Fe ions in the SnO₂lattice as evidenced from the XPS data. However, above 600°, the latticeexpands, approaching the undoped SnO₂ range, and this may be due to theexpulsion of some of the doped Fe ions out of the SnO₂ lattice. Thissuggests that low preparation temperatures provide a relatively uniformdistribution of the Fe dopant ions in the host SnO₂ lattice, and thisfavors ferromagnetism. The high surface diffusion of the dopant atoms insamples annealed above 600° C. causes the gradual disappearance of thisferromagnetic behavior.

4. Temperature Dependence in Sn_(1-x)Co_(x)O₂

To further investigate the role of synthesis parameters on theferromagnetic behavior of 1% Co doped SnO₂, new samples were prepared byannealing the precipitate at temperatures of 250, 350, 450, 600 and 830°C. in air, taking a fresh portion of the dried precipitate each time.The sample annealed at 830° C. showed only the cassiterite phase ofSnO₂, but those annealed at 600, 450, and 350° C. showed cassiterite andorthorhombic phases, as shown in FIG. 5(f). However, the sample annealedat 250° C. showed very strong peaks due to SnO and much weaker peaks ofcassiterite SnO₂. This indicates that lower annealing temperaturespresent a much weaker oxidizing environment, leading to the formation ofSnO and presumably oxygen deficient SnO₂. Some support for thispossibility was obtained from the XRD pattern of a fresh driedprecipitate annealed in flowing H₂ at 350° C., FIG. 5(f), which looksvery similar to the sample prepared in air at 250° C. Based on thisresult, it is inferred that the lower annealing temperatures reduce theoxygen stoichiometry of SnO₂ and eventually converts to SnO at annealingtemperatures <350° C.

In the main panel of FIG. 30, the hysteresis loop measured from the 1%Co doped SnO₂ prepared by annealing the dried precipitate at 350° C. isshown. The saturation magnetization is close to that observed for the˜600° C. annealed samples, but the loop now exhibits a very largecoercivity of 630 Oe with good remenance and squareness. Theseobservations were verified on several batches of samples prepared underidentical conditions. Such large coercivities and squareness of the loopare seldom observed in DMS's at room temperature. Similar reproduciblehysteresis loops were also observed in samples annealed at 450° C.However, no substantial ferromagnetic behavior was detected in samplesannealed in air at 250° C. or in H₂ at 350° C. presumably due toextensive oxygen loss and transformation to the SnO phase. TEM dataobtained from the Sn_(0.99)Co_(0.01)O₂ sample prepared at 350° C., FIG.5(e), showed large micrometer sized particles compared to the ˜37 nmsized particles, FIG. 5(c), observed in samples prepared at 600° C.Since the magnetizations of the samples prepared at these twotemperatures are comparable, as shown in FIG. 30, the observed magneticproperties may not be due to nanoscale size effect. The fact that thesamples prepared by annealing at 350 and 450° C. showed strongferromagnetism with excellent reproducibility—as compared to the lower(˜20%) reproducibility observed when annealed at 600° C. and itscomplete absence in samples annealed at 830° C.—may be closely linked tothe differences in oxygen stoichiometry of these samples. Theoretically,carrier-mediated ferromagnetism in n-type oxide semiconductors isstrongly related to the concentration of oxygen vacancies.

In conclusion, it has been shown that powder samples of chemicallysynthesized Sn_(1-x)CO_(x)O₂ powders with x≦0.01 exhibit RTFM. Thesesamples exhibit significantly high coercivity (˜630 Oe) and goodsquareness of the loop, but with low magnetic moment of 0.133 μ_(B)/Coion. Based on the XRD, PIXE, TEM and magnetic data, the observedferromagnetic interactions seem to be controlled by the oxygenstoichiometry.

II. The Gas Sensing Process

The inventor has developed the ability to detect a gas by causing thegas to flow across a material and measuring the change in a magneticproperty, preferably magnetization, of the material. Changes in themagnetic properties of a material by flowing a gas has never been usedas a sensing method. The preferred material for the process isSn_(0.95)Fe_(0.05)O₂, the manufacture and properties of which have beendescribed above. However, other magnetic materials could be used, solong as their magnetic properties change as a gas flows across them. Agas detected using this process is molecular oxygen, O₂. However, anygas capable of oxidizing or reducing the magnetic material could beused. The preferred apparatus for detecting a gas using this method isshown in FIG. 34. The gas sensor 10 is made from taking anindustry-standard vibrating sample magnetometer (VSM) and adding a gasinlet 20 and mass flow controller (0-300 mL/minute) 22. The VSM and gassensor 10 comprise a VSM controller and power supply 12, a pair ofelectromagnets 14 of +10 kOe, a pair of pickup coils 16, and a VSM headdrive 18. Other combinations of components could be used so long as theyare able to detect magnetic properties, or changes in magneticproperties, of a material. The gas-sensing material 30, preferablySn_(0.95)Fe_(0.05)O₂, is placed at the end of the vibrating sample rod19 connected to the VSM head drive. The gas flows out of the gas inlet20 into a heater (25-600° C.) 21, then passes by the gas sensingmaterial 30, and escapes into the atmosphere. The pickup coils 16measure the magnetization of the gas sensing material 30; themagnetization changes as a function of the flow rate.

To make the gas sensor 10 usable to detect unknown gases, it must firstbe calibrated with known gases. The saturation magnetization ofSn_(0.95)Fe_(0.05)O₂ is shown in FIG. 35 as a function of the flow rateof molecular oxygen. Then, an unknown quantity of gas can flow throughthe gas sensor 10, and the measured magnetization compared to a graphsimilar to that in FIG. 35 to determine the quantity and type of gasflowing through the sensor.

The magnetic properties of the gas sensing material 30 change becausethe carrier-mediated ferromagnetism of Sn_(0.95)Fe_(0.05)O₂ can betailored by exposing it to reducing or oxidizing gaseous atmospheres.Thus, carrier-mediated ferromagnetism has been developed as a new,efficient gas sensing parameter.

Compared to their semiconductor gas sensor counterparts which measurechanges in electrical properties, a magnetic gas sensor is much moreattractive because no electrical contacts are required to detect theresponse, the detection process requires only a moderate magnetic fieldto magnetize the sample and a pickup coil to collect the magneticresponse of the material, powder samples when used offer a very largesurface area and higher sensitivity, magnetic responses are much fasterthan electrical responses, the lack of electrical contacts and the highmagnetic response due to ferromagnetism will further add to thesensitivity of the gas-sensor device, and the operation range can be ashigh as the Curie temperature T_(c).

Since the oxygen stoichiometry in SnO₂, the sensing material (beforedoping), is a surface driven property, the gas-sensing and magneticproperties of a doped oxide semiconductor, when used as the material,are expected to vary significantly with crystalline size, and thereforedepend on the doping concentrations and preparation temperatures.

Although this invention has been described above with reference toparticular means, materials and embodiments, it is to be understood thatthe invention is not limited to these disclosed particulars, but extendsinstead to all equivalents within the scope of the following claims.

1. A method of detecting a gas comprising: causing the gas to flowacross a material; and measuring a change in a magnetic property of thematerial.
 2. The method of claim 1 wherein the material is asemiconductor.
 3. The method of claim 1 wherein the material is atransition metal-doped semiconductor.
 4. The method of claim 1 whereinthe material is a transition metal-doped oxide semiconductor.
 5. Themethod of claim 1 wherein the material is a powder.
 6. The method ofclaim 1 wherein the material is ferromagnetic.
 7. The method of claim 1wherein the material is a powder exhibiting room-temperatureferromagnetism.
 8. The method of claim 1 wherein the material isselected from the group consisting of iron-doped tin dioxide exhibitingroom-temperature ferromagnetism and cobalt-doped tin dioxide exhibitingroom-temperature ferromagnetism.
 9. The method of claim 1 wherein thematerial is iron-doped tin dioxide exhibiting room-temperatureferromagnetism.
 10. The method of claim 1 wherein the material isiron-doped tin dioxide exhibiting room-temperature ferromagnetism. 11.An apparatus for detecting a gas comprising a gas inlet, a flowcontroller, a device which is configured to measure magnetic properties,and a ferromagnetic material.
 12. The apparatus of claim 11 wherein theferromagnetic material is iron-doped tin dioxide.